KR102177591B1 - High-strength steel sheet and its manufacturing method - Google Patents

Abstract

Tensile strength: Provides a high-strength steel sheet having 980 MPa or more and having good formability. Specific component composition, ferrite phase average particle diameter of 1.5 μm or less, ferrite phase area ratio of 2% or more and 15% or less, tempered martensite phase area ratio of 75% or more and 96% or less, per unit area, not tempered It is a high-strength steel sheet having a structure in which the sum of the interface length of the martensite phase and the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/㎡ or more and 5.0 × 10 ¹¹ µm/㎡ or less do.

Description

High-strength steel sheet and its manufacturing method

The present invention relates to a high-strength steel sheet and a manufacturing method thereof.

In recent years, from the viewpoint of preserving the global environment, for the purpose of regulating CO ₂ emissions, improvement in fuel efficiency of automobiles is being pursued throughout the automobile industry. In order to improve fuel efficiency of automobiles, weight reduction of automobiles by thinning of used parts is most effective, and in recent years, the amount of high-strength steel sheets used as materials for automobile parts is increasing.

On the other hand, the formability of the steel sheet tends to deteriorate with increasing strength. Therefore, in addition to high strength, a steel sheet excellent in formability is desired. In the case of a steel sheet having insufficient stretchability, it cannot be applied to automobile parts due to problems such as cracks during molding. In order to reduce the weight of automobile parts and the like, it is essential to develop a steel sheet having both high strength and elongation, and various techniques have been proposed for high strength cold-rolled steel sheets and hot-dip-plated steel sheets that have paid attention to elongation until now.

For example, in Patent Document 1, C: 0.03 to 0.18%, Si: 0.01 to 1.5%, Mn: 0.5 to 3.0%, P: 0.001 to 0.1%, S: 0.0001 to 0.008%, Sol. Al: 0.01 to 0.1%, N: 0.0001 to 0.008%, the remainder is substantially made of Fe, has a steel sheet surface thickness of 10 to 100 µm and a soft layer having a ferrite volume ratio of 90% or more, and the structure of the center It is said that a high-strength cold-rolled steel sheet excellent in workability and corrosion resistance after coating can be obtained by specifying the ratio of the Si concentration in the ferrite phase and in the tempered martensite phase with a tempered martensite volume ratio of 30 to 80%.

In Patent Document 2, by mass%, C: 0.03 to 0.30%, Si: 0.1 to 3.0%, Mn: 0.1 to 5.0%, P: 0.1% or less, S: 0.1% or less, N: 0.01% or less, Al: It contains 0.01 to 1.00%, the balance has a component composition consisting of iron and unavoidable impurities, a tempered martensite having a hardness of more than 380 and not more than 450 Hv contains 70% or more by area ratio, and the remainder is made of ferrite. It is said that by defining the particle size distribution of cementite particles in the tempered martensite, a high-strength cold-rolled steel sheet excellent in balance between elongation and stretch flangeability is obtained.

In Patent Document 3, by mass%, C: 0.05% or more and 0.5% or less, Si: 0.01% or more and 2.5% or less, Mn: 0.5% or more and 3.5% or less, P: 0.003% or more and 0.100% or less, S: 0.02% or less , Al: 0.010% or more and 0.5% or less, B: 0.0002% or more and 0.005% or less, Ti: 0.05% or less, further satisfying Ti> 4N, and the balance being composed of Fe and inevitable impurities, and area It contains 60% or more and 95% or less of tempered martensite by an area ratio, and retained austenite having an average particle diameter of 5% or more and 20% or less by an area ratio of 5 μm or less, and the average particle diameter of the tempered martensite is It is said that a high-strength hot-dip galvanized steel sheet excellent in formability and impact resistance can be obtained by having a structure of 5 μm or less.

In Patent Document 4, as a component composition, in terms of mass%, C: more than 0.10% to less than 0.18%, Si: 0.01 to 1.00%, Mn: 1.5 to 4.0%, P: 0.100% or less, S: 0.020% or less, Al Components containing: 0.010 to 0.500%, Cr: 0.010 to 2.000%, Nb: 0.005 to 0.100%, Ti: 0.005 to 0.100%, B: more than 0.0005% to 0.0030% or less, the remainder consisting of Fe and inevitable impurities It has a composition, and the microstructure is ferrite with an area ratio of 0 to 10%, martensite with an area ratio of 15 to 60%, tempered martensite with an area ratio of 20 to 50%, and bainitic with an area ratio of 20 to 50% It is made of ferrite, and the average grain size of bulk martensite, tempered martensite, and bainitic ferrite is 15 µm or less, respectively, and the value obtained by subtracting the area ratio of tempered martensite from the area ratio of bainitic ferrite is High-strength hot-dip galvanizing with excellent impact resistance and bending workability by specifying the surface layer structure hardness of 20% or less, where the adjacent phase occupied by martensite alone is 5% or less. It is said that a steel plate is obtained.

Japanese Patent Laid-Open Publication No. 2005-256044 Japanese Unexamined Patent Publication No. 2011-52295 Japanese Unexamined Patent Publication No. 2012-31462 Japanese Patent Application Publication No. 2015-117403

However, in the technique proposed in Patent Document 1, there are cases where the strength and bendability are insufficient. Further, in the technique proposed in Patent Document 1, the target structure is obtained by performing two-stage cooling with different cooling rates during annealing, but water cooling is required for cooling in the subsequent stage. Since automobile parts and the like are used in a corrosive environment, it is preferable that the plating properties are good, but in a state where water adheres to the surface of the steel sheet, it cannot be put into a plating bath, and thus Patent Document 1 cannot be applied to a hot-dip plated steel sheet.

In the technology proposed in Patent Document 2, it is considered that it is necessary to control C, which has a significantly faster diffusion rate compared to the substitutional element, and to control the grain growth of cementite in tempered martensite, but control the particle size distribution. It does not disclose how to do so, and it is thought that implementation is impossible.

In the technique proposed in Patent Document 3, the surface layer and the surface layer structure in the center portion are not optimal for bendability. In particular, the retained austenite phase is a structure having high ductility, but in the sheet thickness surface layer subjected to very strict processing, the retained austenite is easy to undergo process-induced transformation, and the martensite phase after transformation may adversely affect the bendability.

In the technique proposed in Patent Document 4, there are cases where Nb is contained and the Nb carbonitride adversely affects the bendability. In addition, bainite ferrite or bainite increases the interface length between the soft bcc structure iron and the hard cementite, so that microcracks resulting from the interface separation between the bcc structure iron and the cementite during bending It is more likely to occur.

The present invention has been made in view of such circumstances, and an object of the present invention is to provide a high-strength steel sheet having a tensile strength of 980 MPa or more and having good formability, and a method for producing the same.

In order to solve the said subject, the present inventors have carefully studied the requirements of a high-strength steel sheet having a tensile strength of 980 MPa or more and excellent formability. Tensile strength: In order to obtain 980 MPa, attention was paid to a tempered martensite phase excellent in both strength and workability. However, even on the tempered martensite, under severe bending conditions, problems such as cracking and local necking were confirmed. Therefore, as a result of pursuing the requirement to improve the bendability, it was discovered that the dislocation was generated by utilizing the transformation deformation accompanying the martensite transformation, or that the stress concentration during bending was disturbed and the deformation was dispersed. . In addition, by presenting a specific amount of a fine ferrite phase having a specific component composition and a specific particle diameter, the area ratio of the tempered martensite phase and the interface length of the untempered martensite phase and the ferrite phase or the tempered martensite phase are determined. When it is in a specific range, the effect of introducing transformational deformation by the non-tempered martensite phase becomes high, and it has become clear that a high-strength steel sheet having a tensile strength of 980 MPa or more and good formability is obtained. In addition, it became clear that the high-strength steel sheet had good plating properties of hot-dip plating.

The present invention has been completed based on the above findings, and its gist is as follows.

[1] By mass%, C: 0.07% or more and 0.14% or less, Si: 0.65% or more and 1.65% or less, Mn: 1.8% or more and 2.6% or less, P: 0.05% or less, S: 0.005% or less, Al: 0.08% Or less, N: 0.0060% or less, Ti: 0.005% or more and 0.030% or less, B: 0.0002% or more and 0.0030% or less, Cr: 0.01% or more and 0.40% or less, and Mo: 0.01% or more and 0.50% or less; or It contains 2 types, satisfies the following formula (1), the remainder has a component composition consisting of Fe and inevitable impurities, the average particle diameter of the ferrite phase is 1.5 µm or less, the area ratio of the ferrite phase is 2% or more and 15% or less, and the temper The area ratio of the de martensite phase is 75% or more and 96% or less, and the total length of the interface length of the untempered martensite phase and the ferrite phase, and the interface length of the untempered martensite phase and the tempered martensite phase per unit area A high-strength steel sheet of 6.3 × 10 ⁸ µm/㎡ or more and 5.0 × 10 ¹¹ µm/㎡ or less.

Here, [%M] (M = Cr, Mo, C) represents the content of each element in mass%, respectively.

[2] In addition to the above component composition, one or two or more selected from V: 0.001% or more and 0.3% or less, Cu: 0.001% or more and 0.1% or less, and Ni: 0.001% or more and 0.2% or less in mass% The high-strength steel sheet according to [1] containing.

[3] The high-strength steel sheet according to [1] or [2], which has a hot-dip plated layer on its surface.

[4] The hot-dip plated layer, by mass%, Fe: 5.0% or more and 20.0% or less, Al: 0.001% or more and 1.0% or less, and Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca , Cu, Li, Ti, Be, Bi, and REM containing one or two or more selected from 0% or more and 3.5% or less in total, and the remainder is the high strength described in [3], having a component composition consisting of Zn and unavoidable impurities. Grater.

[5] The high-strength steel sheet according to [3] or [4], wherein the hot-dip plated layer is an alloyed hot-dip plated layer.

[6] The steel material having the component composition described in [1] or [2] is heated at 1100°C or higher and 1300°C or lower, hot-rolled with a finish rolling of 820°C or higher, and cooling is performed within 3 seconds after the finish rolling. After starting and cooling at an average cooling rate of 30° C./s or more and less than 80° C./s from the finish rolling temperature to 700° C., it is cooled at an average cooling rate of 10° C./s or less to the coiling temperature, and the coiling temperature is 580° C. or more and 680° C. or less. after the hot rolling step, the hot rolling process for winding in, the cold rolling step of performing cold rolling after the cold rolling, more than 500 ℃ (Ac ₃ - 120) with an average heating rate in a temperature range of less than ℃ 4.5 ℃ / s or more to (Ac ₃ - 50) ℃ than (Ac ₃ - 10) heating ℃ to below, and (Ms - 150) ℃ after cooling down to below, in the temperature range of less than 200 ℃ 440 ℃ 15 seconds or more A method of manufacturing a high-strength steel sheet having an annealing process to retain it.

[7] The method for producing a high-strength steel sheet according to [6], which has a hot-dip plating process in which, after the annealing process, it is heated again to 450°C or more and 600°C or less to perform hot-dip plating treatment.

[8] In the hot-dip plating process, the method for producing a high-strength steel sheet according to [7], wherein an alloying treatment is performed after the hot-dip plating treatment.

In addition, in this invention, "high strength" means tensile strength (TS): 980 MPa or more. In addition, the "high-strength steel sheet" is a cold-rolled steel sheet or a hot-dip plated steel sheet having a hot-dip plated layer on the surface of the cold-rolled steel sheet. The "hot-dip plated steel plate" includes not only a hot-dip plated steel plate but also an alloyed hot-dip plated steel plate.

The high-strength steel sheet of the present invention has both tensile strength (TS): high strength of 980 MPa or more and excellent formability. Therefore, when the high-strength steel sheet of the present invention is applied to automobile parts, especially automobile skeleton members, additional weight reduction of automobile parts can be realized. Moreover, since the high-strength steel sheet of the present invention is also excellent in the plating properties of hot-dip plating, it is possible to obtain a hot-dip-plated steel sheet having both high strength and excellent formability and suppressing the occurrence of sub-plated portions.

Hereinafter, an embodiment of the present invention will be described.

＜Constituent composition of high strength steel sheet＞

The component composition of the high-strength steel sheet of the present invention is mass%, C: 0.07% or more and 0.14% or less, Si: 0.65% or more and 1.65% or less, Mn: 1.8% or more and 2.6% or less, P: 0.05% or less, S: 0.005 % Or less, Al: 0.08% or less, N: 0.0060% or less, Ti: 0.005% or more and 0.030% or less, B: 0.0002% or more and 0.0030% or less, Cr: 0.01% or more and 0.40% or less, and Mo: 0.01% or more 0.50% It contains 1 type or 2 types selected from the following, and satisfies the following formula (1).

Here, [%M] (M = Cr, Mo, C) represents the content of each element in mass%, respectively.

Each component composition will be described below. In the following description, "%" indicating the content of a component means "mass%".

C: 0.07% or more and 0.14% or less

C relates to the formation of a non-tempered martensite phase and a tempered martensite phase, and the strength thereof. Tensile strength: In order to obtain 980 MPa or more, it is necessary to contain at least 0.07% or more of C content. On the other hand, when the C content exceeds 0.14%, coarse cementite is formed and the bendability decreases. Therefore, in the present invention, the C content is 0.07% or more and 0.14% or less. The lower limit of the C content is preferably 0.08% or more, more preferably 0.09% or more, and still more preferably 0.10% or more. The upper limit of the C content is preferably 0.13% or less, more preferably 0.12% or less, and still more preferably 0.11% or less.

Si: 0.65% or more and 1.65% or less

Si is an element that increases work hardenability and contributes to improvement of bendability. Si is also involved in ferrite phase formation, and the high-strength steel sheet of the present invention is annealed by rapid heating at the time of manufacture, but if Si is less than 0.65%, the area ratio of the desired ferrite phase is not stably obtained, resulting in lower bendability. . From the above, in order to obtain the bendability required by the present invention, it is necessary to contain at least 0.65% or more of Si. Moreover, when Si exceeds 1.65%, the adverse influence on the plating property becomes present. Therefore, in the present invention, the Si content is 1.65% or less. With respect to the lower limit of the Si content, it is preferably 0.80% or more, more preferably 0.90% or more, and still more preferably 1.00% or more. The upper limit of the Si content is preferably 1.60% or less, more preferably 1.50% or less, and still more preferably 1.40% or less.

Mn: 1.8% or more and 2.6% or less

Mn is an element that improves etchability and contributes to suppression of formation of a coarse ferrite phase. In order to obtain the quenching property required in the present invention, it is necessary to contain 1.8% or more of Mn. On the other hand, since Mn lowers the Ms point, if the amount of Mn is too large, the untempered martensite phase specified in the present invention cannot be obtained. Further, since Mn lowers the martensite transformation initiation temperature, in the annealing process when manufacturing the high-strength steel sheet of the present invention, a very high cooling rate in cooling to a cooling stop temperature (Ms-150°C) or lower It becomes necessary. From the viewpoint of deteriorating the controllability of such a cooling rate, the bendability deteriorates. Moreover, Mn deteriorates the plating property. From the above, in the present invention, the Mn content is 1.8% or more and 2.6% or less. The lower limit of the Mn content is preferably 1.9% or more, more preferably 2.0% or more, and still more preferably 2.1% or more. The upper limit of the Mn content is preferably 2.5% or less, more preferably 2.4% or less, and still more preferably 2.3% or less.

P: 0.05% or less

P is an element that segregates at grain boundaries and deteriorates bendability. In the present invention, the P content can be allowed up to 0.05% or less. P content is 0.05% or less, preferably 0.04% or less. Although it is preferable to reduce the P content as much as possible, 0.002% may inevitably be mixed in production.

S: 0.005% or less

S forms coarse MnS in the steel, which extends at the time of hot rolling and becomes wedge-shaped inclusions, thereby causing adverse effects on bendability. In the present invention, the S content can be allowed to 0.005%. The S content is 0.005% or less, preferably 0.003% or less. Although it is preferable to reduce the S content as much as possible, 0.0002% may inevitably be mixed in production.

Al: 0.08% or less

When Al is added as a deoxidizing agent in the step of steel making, it is preferable to contain Al content of 0.02% or more. More preferably, it is 0.03% or more. On the other hand, Al forms an oxide that deteriorates the bendability. Therefore, in the present invention, the Al content is 0.08% or less, preferably 0.07% or less. More preferably, it is 0.06% or less, More preferably, it is 0.05% or less.

N: 0.0060% or less

N is a harmful element that adversely affects the bendability by bonding with Ti or B to form coarse inclusions, lowers the quenching property, and hinders the formation of a fine ferrite phase. In the present invention, the N content can be allowed to 0.0060%. The N content is preferably 0.0050% or less. More preferably, it is 0.0040% or less. Although it is preferable to reduce the N content as much as possible, 0.0005% may inevitably be mixed in production.

Ti: 0.005% or more and 0.030% or less

Ti is an element capable of fixing N, which is a harmful element, as a nitride, and preventing a decrease in etchability of B due to N. In order to suppress the decrease in quenchability due to B, it is necessary to contain at least 0.005% or more of Ti. On the other hand, when Ti exceeds 0.030%, the bendability falls due to the carbonitride containing coarse Ti. Therefore, in the present invention, the Ti content is 0.005% or more and 0.030% or less. In addition, it is preferable to limit the Ti content to 0.010% or more and 0.028% or less, and to limit Nb, which is an element capable of forming other coarse inclusions, to less than 0.003%.

B: 0.0002% or more and 0.0030% or less

B is an element that improves the quenching property and contributes to the formation of a fine ferrite phase. In order to obtain this effect, it is necessary to contain at least 0.0002% or more of B. On the other hand, when the B content exceeds 0.0030%, the bendability decreases due to the adverse effect of the ductility decrease caused by the solid solution B. Therefore, in the present invention, the B content is 0.0002% or more and 0.0030% or less. With respect to the lower limit of the B content, it is preferably 0.0005% or more, more preferably 0.0006% or more, and still more preferably 0.0010% or more. With respect to the upper limit of the B content, it is preferably 0.0025% or less, more preferably 0.0020% or less, and still more preferably 0.0016% or less.

Cr: 0.01% or more and 0.40% or less and Mo: 0.01% or more and 0.50% or less.

Cr and Mo are elements which improve quenching property, and are elements that contribute to formation of a fine ferrite phase. In order to obtain this effect, it is necessary to contain Cr by 0.01% or more, Mo to contain 0.01% or more, or to contain Cr and Mo by 0.01% or more, respectively. On the other hand, when Cr exceeds 0.40%, the above effect is saturated, so the upper limit is set to 0.40%. In addition, if the Cr content exceeds 0.40%, the plating property is lowered, and a steel sheet having good plating properties is difficult to obtain. Therefore, it is preferable to be 0.40% or less from the viewpoint of plating properties. On the other hand, when the Mo content exceeds 0.50%, the transformation point will be out of the appropriate range, and the untempered martensite phase specified in the present invention will not be obtained. Therefore, in the present invention, the Cr content is 0.40% or less, and the Mo content is 0.50% or less. The lower limit of the Cr content is preferably 0.02% or more, more preferably 0.03% or more, and still more preferably 0.04% or more. The upper limit of the Cr content is preferably 0.35% or less, more preferably 0.30% or less, and still more preferably 0.20% or less. With respect to the lower limit of the Mo content, it is preferably 0.02% or more, more preferably 0.05% or more, and still more preferably 0.10% or more. The upper limit of the Mo content is preferably 0.43% or less, more preferably 0.40% or less, and still more preferably 0.30% or less.

(1) satisfy the expression

In addition to suppression of formation of a coarse ferrite phase by improvement of quenchability during annealing, Cr and Mo interfere with the austenite-to-ferrite interfacial movement in the cooling process after finish rolling in the hot rolling process. The effect of making the hot-rolled structure fine is also obtained. By making the hot-rolled structure fine, many portions of the hot-rolled sheet with a high C concentration are produced to obtain a non-tempered martensite structure specified in the present invention. However, when the C content is large, the portion with a high C concentration is The distribution density becomes coarse, and the structure specified in the present invention cannot be obtained. In order to suppress this, it is necessary to contain Cr and Mo within a range that satisfies the formula (1). The coefficient of Mo represents the magnitude of the influence of changing the distribution density of C. (1) A preferable range of the left side of the equation is 2.1 or more. (1) There is no upper limit of the formula, and it is determined by the upper limit of the content of Cr and Mo.

Although the above is the basic configuration of the component composition of the high-strength steel sheet of the present invention, the high-strength steel sheet of the present invention is further, in mass%, V: 0.001% or more and 0.3% or less, Cu: 0.001% or more and 0.1% or less, and Ni: 0.001 You may contain 1 type or 2 or more types selected from% or more and 0.2% or less.

V, Cu, and Ni are elements contributing to further strengthening. By containing these, the strength stability is improved. In order to obtain this effect, it is preferable to contain at least 0.001% or more of V, Cu, and Ni, respectively. On the other hand, when the V content is more than 0.3%, the Cu content is 0.1%, and the Ni content is more than 0.2%, the transformation point of the steel sheet changes, and the structure required by the present invention is difficult to obtain. The preferred V content is 0.01% or more and 0.2% or less, the preferred Cu content is 0.01% or more and 0.08% or less, and the preferred Ni content is 0.01% or more and 0.1% or less.

Components other than the above components are Fe and unavoidable impurities. In addition, in the above component composition, when a component that does not need to be contained is contained below the lower limit, the component is considered to be contained as an unavoidable impurity.

Moreover, it is preferable that the high-strength steel sheet of the present invention further satisfies the following formula (2).

[% Cr] ≤ 0.215 [% Si] 2 - 0.8 [% Si] + 0.747 (2)

Here, [%M] (M = Cr, Si) represents the content of each element in mass%, respectively.

(2) satisfies the expression

When both Si and Cr are contained, plating properties deteriorate due to a synergistic effect. Therefore, from the viewpoint of plating properties, the upper limit of the Cr content changes with respect to the Si content. Equation (2) is an approximation of the relationship between the upper limit of Cr at which the plating property becomes good when Si is changed. (2) As long as it is a range that satisfies the formula, a steel sheet having good plating properties can be obtained.

<Metal structure of high-strength steel plate>

Subsequently, the metal structure (steel structure) of the high-strength steel sheet of the present invention will be described. The metal structure is obtained by observing the structure of a high-strength steel sheet parallel to the rolling direction using a scanning electron microscope.

The metal structure of the high-strength steel sheet of the present invention has a ferrite phase average particle diameter of 1.5 μm or less, a ferrite phase area ratio of 2% or more and 15% or less, a tempered martensite phase area ratio of 75% or more and 96% or less, per unit area. , The total length of the interface length of the non-tempered martensite phase and the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/㎡ or more and 5.0 × 10 ¹¹ µm/㎡ or less. .

The average particle diameter of the ferrite phase is 1.5 µm or less, and the area ratio of the ferrite phase is 2% or more and 15% or less

The ferrite phase is a structure having ductility and has an effect of improving bendability. If there is no ferrite phase at all, work hardenability and ductility are insufficient, and cracks occur due to insufficient ductility or stress concentration during bending. On the other hand, when the ferrite grains are coarse, the formation of a finely tempered martensite phase is inhibited, and the bendability is rather deteriorated. In addition, when there are too many ferrite phases, the tensile strength of 980 MPa or more is difficult to obtain in the non-tempered martensite phase and the tempered martensite phase in the high-strength steel sheet of the present invention. Therefore, from the viewpoint of strength and bendability, it is necessary to control the average particle diameter and area ratio of the ferrite phase together. In the present invention, the average particle diameter of the ferrite phase is 1.5 μm or less, and the area ratio of the ferrite phase is 2% or more and 15% or less. Needs to be. Preferably, the average particle diameter of the ferrite phase is 1.2 µm or less, and the area ratio of the ferrite phase is 2% or more and 10% or less. The lower limit of the average particle diameter of the ferrite phase is not particularly limited, but, for example, the average particle diameter of the ferrite phase is 0.1 µm or more.

The area ratio of the tempered martensite phase is 75% or more and 96% or less

The tempered martensite phase is a structure in which fine iron-based carbides having orientation in crystal grains and corrosion marks are observed. Examples of iron-based carbides include cementite, η carbide, χ carbide, and ε carbide. This tempered martensite phase is excellent in balance between strength and ductility, and in the present invention, strength is obtained mainly from the tempered martensite phase. If the area ratio of the tempered martensite phase is less than 75%, the tensile strength tends to decrease, and although it may vary depending on the average particle diameter of the ferrite phase and the area ratio of the ferrite phase, the area ratio of the tempered martensite phase is 75%. If it is less than, the tensile strength may be less than 980 MPa. On the other hand, when the area ratio of the tempered martensite phase exceeds 96%, the good bendability required by the present invention cannot be obtained. From the above, in the present invention, the area ratio of the tempered martensite phase is 75% or more and 96% or less. It is more preferable that the area ratio of the tempered martensite phase is more than 85% from the point in which the variation in tensile properties tends to decrease as the structure becomes uniform. More preferably, it is 86% or more. The upper limit is preferably 94% or less, and more preferably 91% or less. In the tempered martensite phase, for example, the martensite phase produced by rapid cooling after heating in the annealing step in the manufacturing method of the high-strength steel sheet of the present invention described later is in a temperature range of 200°C or more and 440°C or less. It can be caused by changes in the process of staying (maintaining) in.

Per unit area, the total length of the interface length of the non-tempered martensite phase and the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/㎡ or more 5.0 × 10 ¹¹ µm/ ㎡ or less

The non-tempered martensite phase is a structure in which no iron-based carbide is found in the mouth, and is observed with a white contrast under a scanning electron microscope. Per unit area, the total length of the interface length of the non-tempered martensite phase and the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/㎡ or more 5.0 × 10 ¹¹ µm If it is less than /m 2, it will be a steel sheet having good bendability. If the interface length is 6.3 × 10 ⁸ µm/m 2 or more and 5.0 × 10 ¹¹ µm/m 2 or less, it is assumed that the stress concentration at the time of bending is prevented by introducing a suitable dislocation into the inside of the steel sheet, and good bendability is obtained. The martensite phase that has not been tempered is obtained in detail by the method for producing a high-strength steel sheet of the present invention described later. Dislocations are introduced into adjacent tissues due to transformational deformation at the time of formation of this non-tempered martensite phase. If the adjacent structure of the non-tempered martensite phase is the non-tempered martensite phase, the effect of improving the bendability by introducing dislocations is not obtained. Therefore, it is necessary to have a region adjacent to the ferrite phase or the tempered martensite phase on the non-tempered martensite phase, and in order to obtain an effect of improving the bendability, the non-tempered martensite phase per unit area and The total value of the interface length of the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase must be 6.3×10 ⁸ µm/m 2 or more. On the other hand, if the interface length exceeds 5.0 × 10 ¹¹ µm/m 2, it may be because dislocations are excessively introduced, or the bendability decreases. The total value of the interface length of the non-tempered martensite phase and the ferrite phase and the interface length of the non-tempered martensite phase and the tempered martensite phase per unit area is preferably 8.0 × 10 ⁸ µm/m2 with respect to the lower limit. Or more, more preferably 1.0×10 ¹⁰ μm/m 2 or more. The upper limit is preferably 4.6 × 10 ¹¹ μm/m 2 or less. More preferably, it is 20 × 10 ¹⁰ μm/m 2 or less.

In addition to the ferrite phase, the tempered martensite phase, and the untempered martensite phase, the residual austenite phase (residue γ) and the bainite phase (B) may be contained, for example, about 4% or less. Further, in the present invention, if the interface length is within a specific range, the area ratio of martensite that has not been tempered is not particularly limited, but it is often 1 to 5%.

<Molten plating layer>

The high-strength steel sheet of the present invention may be one having a hot-dip plated layer on its surface, that is, a hot-dip plated steel plate or an alloyed hot-dip plated steel plate. The hot-dip plating layer will be described below. In the present invention, the component constituting the hot-dip plating layer is not particularly limited and may be a general component. Examples of the hot-dip plating layer include a Zn-based plating layer and an Al-based plating layer. Examples of the Zn-based plating include general hot dip galvanizing (GI), Zn-Ni-based plating, and Zn-Al-based plating. In addition, Al-Si-based plating (eg, Al-Si-based plating containing 10 to 20 mass% of Si) can be exemplified as Al-based plating. As the Zn-based plating layer, specifically, by mass%, Fe: 5.0 to 20.0%, Al: 0.001% to 1.0%, and further Pb, Sb, Si, Sn, Mg, Mn, Ni , Cr, Co, Ca, Cu, Li, Ti, Be, Bi, REM containing one or two or more selected from 0 to 3.5% in total, and the balance is a hot-dip galvanized layer consisting of Zn and unavoidable impurities. Can be lifted. Moreover, the hot-dip plated layer may be an alloyed hot-dip plated layer. As an alloyed hot dip plating layer, an alloyed hot dip galvanizing (GA) layer is mentioned, for example. Although the adhesion amount of plating is also arbitrary, it is preferable to set it as 120 g/m<2> or less per one side from a weldability viewpoint. In addition, the lower limit of the adhesion amount is not particularly limited, but is usually 30 g/m 2 or more.

The sheet thickness of the high-strength steel sheet of the present invention is not particularly limited, but it is preferably 0.5 mm or more and 2.6 mm or less. When the high-strength steel sheet has a plating layer, the sheet thickness is the sheet thickness of the base steel sheet excluding the plating layer.

<Method of manufacturing high strength steel plate>

Next, a method of manufacturing the high-strength steel sheet of the present invention will be described. The manufacturing method of a high-strength steel sheet of this invention has a hot rolling process, a cold rolling process, and an annealing process. Specifically, in the method of manufacturing a high-strength steel sheet of the present invention, a steel material having the above component composition is heated at 1100°C or higher and 1300°C or lower, and hot rolling is performed at a finish rolling of 820°C or higher, and 3 seconds after finish rolling. After starting cooling within the range from the finish rolling temperature to 700°C and cooling at an average cooling rate of 30°C/s or more and less than 80°C/s, cooling to the coiling temperature at an average cooling rate of 10°C/s or less, and a coiling temperature of 580°C or more after the hot rolling step, the hot rolling process for winding below 680 ℃, and the cold rolling step of performing cold rolling after the cold rolling, more than _{500 ℃ (Ac 3 - 120)} ℃ average heating of the temperature range of less than the speed is more than _{4.5 ℃ / s (Ac 3 -} 50) ℃ than (Ac ₃ - 10) heating ℃ to the following, and - after cooling to (Ms 150) ℃ or less, in a temperature range of less than 200 ℃ 440 ℃ It has an annealing process of staying longer than 15 seconds. In the case of manufacturing a high-strength steel sheet having a hot-dip plated layer on its surface, a hot-dip plating process is provided in which, after the annealing process, it is heated again to 450° C. or more and 600° C. or less to perform hot-dip plating treatment. In addition, in the case of manufacturing a high-strength steel sheet having an alloyed hot-dip plated layer on its surface, in the hot-dip plating process, an alloying treatment is performed after the hot-dip plating treatment. Hereinafter, each process will be described in detail. In addition, in the following description, the temperature is taken as the surface temperature of a steel material or a steel plate unless otherwise specified. In addition, the average heating rate is ((surface temperature after heating-surface temperature before heating)/heating time), and the average cooling rate is ((surface temperature before cooling-surface temperature after cooling)/cooling time).

(Hot rolling process)

In the hot rolling process, the steel material having the above component composition is heated at 1100° C. or higher and 1300° C. or lower, hot rolling is performed at a finish rolling of 820° C. or higher, and cooling is started within 3 seconds after the finish rolling, and from the finish rolling temperature. This is a step of cooling at an average cooling rate of 30°C/s or more and less than 80°C/s to 700°C, cooling to a winding temperature at an average cooling rate of 10°C/s or less, and winding up at a winding temperature of 580°C or more and 680°C or less.

The solvent method for manufacturing the steel material is not particularly limited, and a known solvent method such as a converter or an electric furnace may be employed. Moreover, you may perform secondary refining in a vacuum degassing furnace. After that, it is preferable to make a slab (steel material) by a continuous casting method from the problem of productivity and quality. Moreover, it is good also as a slab by a well-known casting method, such as a coarse-disintegration rolling method and a thin slab playing method.

Steel material heating temperature: 1100 ℃ or more and 1300 ℃ or less

In the present invention, it is necessary to heat the steel material prior to rough rolling to make the steel structure of the steel material a substantially homogeneous austenite phase. Moreover, in order to suppress the formation of coarse inclusions, control of the heating temperature becomes important. If the heating temperature is less than 1100°C, the desired finished rolling temperature cannot be obtained. On the other hand, when the heating temperature exceeds 1300°C, scale loss increases, and damage to the furnace body of the heating furnace increases. Therefore, it is necessary to set the heating temperature of the steel material to 1100°C or more and 1300°C or less. With respect to the lower limit of the heating temperature of the steel material, it is preferably 1120°C or higher, and more preferably 1150°C or higher. The upper limit is preferably 1280°C or less, more preferably 1260°C or less. In addition, the rough rolling conditions of the rough rolling after heating are not particularly limited.

Finish rolling temperature: 820 ℃ or more

Finish rolling after rough rolling is performed at a temperature of 820°C or higher. When the finish rolling temperature is less than 820°C, austenite-to-ferrite transformation starts during rolling, and ferrite grains grow grains, so that the steel sheet structure specified in the present invention cannot be obtained. Therefore, in the manufacturing method of this invention, the finish rolling temperature is 820 degreeC or more. The finish rolling temperature is preferably 840°C or higher. The upper limit of the finish rolling temperature is not particularly limited, but is usually 1000°C or less.

After finishing the finish rolling, cooling was started within 3 seconds, the average cooling rate was 30°C/s or more and less than 80°C/s from the finish rolling temperature to 700°C, and then cooled to the winding temperature at an average cooling rate of 10°C/s or less. , Winding at a winding temperature of 580°C or more and 680°C or less.

In the hot-rolled steel sheet, in order to increase the interface length of the martensite phase which is not tempered to be within the specific range, it is necessary to form a structure having a fine ferrite structure. After finishing the finish rolling, if the time from the finish rolling temperature to 700° C. to the start of cooling performed at an average cooling rate of 30° C./s or more exceeds 3 seconds (s), austenite → ferrite transformation starts at high temperature. It is discarded, and a fine ferrite structure is not obtained. From the start of cooling, the austenite-ferrite grain boundary movement speed is similarly fast up to 700° C., which is a temperature range in which ferrite grains tend to coarsen. Therefore, the average cooling rate from the finish rolling temperature to 700° C. needs to be 30° C./s or more. On the other hand, when the cooling rate is cooled to 80° C./s or more, the control of the cooling stop temperature becomes unstable, and there is a possibility that C diffusion at 700° C. or less described later may not be promoted, and a desired steel sheet structure cannot be easily obtained. At 700° C. or less, it is necessary to diffuse C at the triple points of the ferrite interface and grain boundary to form a high-concentration portion, and to promote the formation of an untempered martensite phase. Therefore, the average cooling rate from 700°C to the coiling temperature needs to be 10°C/s or less. After finishing the finish rolling, preferably within 0.1 seconds or more and 2.0 seconds, the average cooling rate from the finish rolling temperature to 700° C. is cooled at an average cooling rate of 35° C. or more and less than 80° C./s, and then cooled to a coiling temperature at an average cooling rate of 8° C./s or less. Do. In addition, cooling performed at an average cooling rate of 30°C/s or more and less than 80°C/s from the finish rolling temperature to 700°C, and cooling performed at an average cooling rate of 10°C/s or less to the subsequent coiling temperature may be performed continuously. However, after cooling performed at an average cooling rate of 30° C./s or more and less than 80° C./s from the finish rolling temperature to 700° C., from the viewpoint of the distribution of the C concentration, 3 seconds or more and 15 seconds or less at 680° C. or more and 700° C. or less It is preferable to stay without forced cooling during (i.e., without performing a special cooling operation), and then cool to the coiling temperature at an average cooling rate of 10° C./s or less. In addition, the lower limit of the average cooling rate up to the coiling temperature is not particularly limited, but is usually 1° C./s or more.

The coiling temperature is 580 ℃ or more and 680 ℃ or less

In winding, it becomes important not to change the generated fine ferrite structure and not to generate a low-temperature transformed phase such as a bainite phase. If the temperature is lower than 580°C, there is a possibility that bainite transformation may be initiated, and a martensite structure that is not tempered as defined in the present invention cannot be obtained. On the other hand, when the coiling temperature exceeds 680°C, grain growth of ferrite grains reduces the interfacial length of the non-tempered martensite phase, the ferrite phase, and the tempered martensite phase. From the above, it is necessary that the range of the coiling temperature be 580°C or more and 680°C or less. The lower limit is preferably 600°C or higher, and more preferably 610°C or higher. The upper limit is preferably 660°C or less, and more preferably 650°C or less.

(Cold rolling process)

The cold rolling process is a process of cold rolling a hot-rolled sheet after the hot rolling process. In order to obtain a desired sheet thickness, it is necessary to perform cold rolling on the hot rolled sheet after the hot rolling step. In the present invention, the conditions of the cold rolling process are not particularly limited. From the viewpoint of the plate shape at the time of cold rolling, it is preferable to make the rolling rate of cold rolling into 40 to 80%.

(Annealing process)

Annealing process after cold rolling, at least 500 ℃ (Ac ₃ - 120) with an average heating rate is more than 4.5 ℃ / s in the temperature range of less than ℃ (Ac ₃ - 50) ℃ than (Ac ₃ - 10) ℃ below It is a step of heating to and cooling to (Ms-150)°C or lower, and then allowing it to remain in a temperature range of 200°C or higher and 440°C or lower for 15 seconds or longer. The annealing process is performed, for example, in an oxidation-free or direct-fired continuous annealing line, and when forming a hot-dip plating layer, it is performed in an oxidation-free or direct-fired continuous hot-dip plating line.

More than _{500 ℃ (Ac 3 - 120)} ℃ average heating rate of 4.5 ℃ / s or more in the temperature range of less than

In the hot rolling step, a fine ferrite phase was used to obtain the martensite phase specified in the present invention, but the effect of this fine ferrite phase was lost when the recovery proceeds during heating in the annealing step, thereby inhibiting recovery. It is necessary to proceed with recrystallization or ferrite → austenite transformation. Since the - (120 Ac ₃₎ ℃ below the recrystallization, or ferrite → austenite transformation starts at the, in the present invention, after the cold rolling process recovery is in progress and, in more than 500 ℃ - than _{(Ac 3 50) ℃ (Ac} 3 When heating to -10)°C or lower, the average heating rate in a temperature range of 500°C or higher (Ac ₃ -120)°C or lower needs to be 4.5°C/s or higher. Ac ₃ is a temperature at which ferrite completes transformation into austenite upon heating. The average heating rate in a temperature range of 500°C or higher (Ac ₃ -120)°C or lower is preferably 5.0°C/s or higher. Further, the cold rolling process after (Ac ₃ - 50) ℃ than (Ac ₃ - 10) when hot-℃ to less than, less than 500 ℃ or, - a heating rate of from higher than (Ac ₃ 120) ℃ temperature is not particularly limited It doesn't work. In addition, the upper limit of the average heating rate is not particularly limited, but is usually 50° C./s or less.

The heating temperature (annealing _{temperature): (Ac 3 - 50)} ℃ than (Ac ₃ - 10) ℃ below

In the heating step, it is necessary to form an appropriate amount of the ferrite phase. If the heating temperature is outside the range of (Ac ₃ -50)°C or higher (Ac ₃ -10)°C or lower, the area ratio of the ferrite phase specified in the present invention cannot be obtained. The heating temperature is preferably a - - (15 Ac ₃₎ ℃ than (Ac ₃ 40) ℃ above. In addition, (Ac ₃ - 50) ℃ than (Ac ₃ - 10) ℃ residence time in the following is not particularly limited, for example, is not more than 300 seconds, preferably for more than 30 seconds to 300 seconds, and more preferably It is more than 50 seconds and less than 250 seconds.

Cooling stop temperature after heating (cooling end temperature): (Ms-150) ℃ or less

In cooling after heating, it is necessary to make most of the structure into a martensite shape. At this time, if the martensite phase is not formed but remains as austenite, a coarse untempered martensite phase or retained austenite is finally formed, thereby reducing the bendability. To avoid the adverse effect, (Ac ₃ - 50) ℃ than (Ac ₃ - 10) is a cooling-stop temperature after the heating at below ℃ (Ms - 150) is required to be less than ℃, i.e., (Ac ₃ - 50 ) ℃ than (Ac ₃ - it is necessary to cool to 150) ℃ below - 10) after the heating at below ℃ (Ms. In addition, Ms is a temperature at which austenite begins to transform into a martensite phase during cooling. In order to stably obtain a desired structure, the cooling stop temperature after heating is preferably made at (Ms-170)°C or lower (Ms-300)°C or higher. Also, it helps in the present invention, the rise ?? chingseong by Cr, Mo and _{B, (Ac 3 - 50)} ℃ than (Ac ₃ - 10) 550 that is a fear of the ferrite grain growth on from the hot end of the below ℃ If the average cooling rate up to °C is less than 30 °C/s, there is a risk of grain growth of ferrite grains, so the average cooling rate up to 550 °C is preferably 35 °C/s or more, and 40 °C/s or more. It is more preferable to do it. The upper limit of the average cooling rate is not particularly limited, but is usually 70° C./s or less. In addition, after cooling to (Ms-150)°C or lower, the process is kept at a temperature of (Ms-150)°C or lower, for example, 7 s or more and 50 s or less, and then stays at 200°C or more and 440°C or less for 15 s or more It is preferable to carry out.

Stays for 15 s or more at 200 ℃ or more and 440 ℃ or less

(Ms-150) In order to make the martensite phase obtained by cooling to C or less to become a tempered martensite phase, it is necessary to stay (maintain) for 15 s or more in a temperature range of 200° C. or more and 440° C. or less. In addition, by staying (maintaining) for 15 s or more in a temperature range of 200°C to 440°C, C is distributed to the austenite phase, and the Ms point is locally lowered to promote the formation of untempered martensite. have. If the cooling stop temperature is less than 200°C, it is heated to a temperature range of 200°C or more and 440°C or less. If the residence temperature (maintenance temperature) in the temperature range of 200°C or more and 440°C or less is less than 200°C, the diffusion of carbon contained in the martensite phase by supersaturation is slow, so that a sufficient tempering effect cannot be obtained and the bendability deteriorates. If it stays at a temperature exceeding 440° C., the martensite phase is excessively tempered, and austenite is decomposed to prevent untempered martensite from being obtained, so that the high strength specified in the present invention cannot be obtained. In addition, if the residence time is less than 15 s, a sufficiently tempered martensite phase cannot be obtained, so that the bendability decreases. Preferred residence conditions are to be kept for 20 s or more at 250°C or more and 430°C or less. In addition, the upper limit of the residence time is not particularly limited, but is usually 100 seconds or less.

(Melt plating process)

The hot-dip plating process is a process of hot-dip plating treatment by heating again to 450°C or more and 600°C or less after the annealing process. Thereby, a high-strength steel sheet having a hot-dip plated layer, that is, a hot-dip plated steel sheet is obtained. Further, in the hot-dip plating process, by further alloying treatment after hot-dip plating treatment, a high-strength steel sheet having an alloyed hot-dip plated layer, that is, an alloyed hot-dip plated steel sheet, is obtained.

Reheat to 450 ℃ or more and 600 ℃ or less

In order to obtain a hot-dip plated steel sheet, it is necessary to immerse the steel sheet after the annealing process in a plating bath. The component composition of the plating bath may be the same as that of the hot-dip plating layer to be produced. From the viewpoint of the appearance quality of the hot-dip plated steel sheet, the reheating temperature needs to be 450°C or higher. On the other hand, when performing the alloying treatment, if the heating temperature is excessively increased, the strength decreases and the desired tensile strength cannot be obtained. In the present invention, up to 600°C can be allowed. Therefore, in this invention, the reheating temperature is 450 degreeC or more and 600 degreeC or less. The temperature of the plating bath is preferably 450°C or more and less than 500°C. Further, the alloying treatment temperature is preferably 500°C or more and 600°C or less.

In addition, in the case of manufacturing a cold-rolled steel sheet having no plating layer, it is preferable to quench it to room temperature using water or the like after staying for 15 seconds or more in a temperature range of 200°C or more and 440°C or less in the annealing step. Further, in the case of manufacturing a cold-rolled steel sheet having a plating layer on its surface, it is preferable to perform a hot-dip plating step of hot-dip plating by heating again to 450°C or more and 600°C or less, and then rapid cooling to room temperature using water or the like. Room temperature is 0°C or more and 50°C or less. Moreover, rapid cooling refers to cooling of 20 degreeC/s or more of cooling rates.

Example

A steel material with a thickness of 250 mm having the component composition shown in Table 1 and the remainder consisting of Fe and unavoidable impurities was subjected to a hot rolling process under the hot rolling conditions shown in Table 2 to obtain a hot rolled sheet, and the cold rolling rate was 40% or more and 80% The following cold rolling process was performed, and it was set as the cold-rolled plate of 1.0-2.0 mm plate|board thickness, and the annealing process of the conditions shown in Table 2 was performed, and the steel plate was obtained. Thereafter, the obtained steel sheet was subjected to a hot-dip plating treatment to form a hot-dip galvanized layer on the surface (GI material). Further, for some of the steel sheets in which the hot-dip galvanized layer was formed, after forming the hot-dip galvanized layer, alloying treatment was performed at the alloying temperature shown in Table 2 to form an alloyed hot-dip galvanized layer (GA material). The steel sheet up to the annealing process was produced in a direct-heating type continuous annealing line, and the hot-dip plated layer or an alloyed hot-dip plated layer was formed on a direct-heating type continuous hot-dip plating line. Here, the temperature of the plating bath (plating composition: Zn-0.13 mass% Al) immersed in the direct-fired continuous hot-dip plating line is 460°C, and the amount of plating deposited is GI material (melt plated steel sheet), GA material (alloyed hot-dip plated steel sheet ) All were 45 to 65 g/m 2 per side, and the amount of Fe contained in the plating layer was in the range of 6 to 14 mass%. Ac ₃ points were obtained from the transformation expansion curve obtained at a heating rate of 6°C/s using a thermal expansion measuring device. In addition, the Ms point was obtained from a transformation curve obtained by using a thermal expansion measuring device to obtain a cooling rate from the Ac ₃ point to 300°C at 30°C/s after heating to an Ac ₃ point or more.

In addition, in the hot rolling step, after cooling performed at an average cooling rate of 30°C/s or more from the finish rolling temperature to 700°C, the “retention time” of the hot rolling step of Table 2 in a temperature range of 680°C to 700°C. It maintained for the time indicated in the column (for 5 to 10 seconds), and then cooled to the coiling temperature at an average cooling rate of 10°C/s or less.

Also, in the annealing step, (Ac ₃ - 50) ℃ than (Ac ₃ - 10) more than 500 ℃ for an average heating rate of the step of heating to below ℃, - only between (Ac ₃ 120) ℃ below Table 2 Although it described in, (Ac ₃ - 50) ℃ than (Ac ₃ - 10) at the time of heating to below ℃, above 500 ℃ described in Table 2 - the same as the average heat rate of between (Ac ₃ 120) ℃ below I did.

Also, in the annealing process, in Table 2 at a temperature not higher than _{- (10 Ac 3) ℃ (} Ac 3 - 50) ℃ than (Ac ₃ - - 10) After heating ℃ to below, at least (Ac ₃ 50) ℃ The time (51 to 169 s) described in the "retention time" column of the annealing step was maintained (retained), and then a step of cooling to (Ms-150)°C or lower was performed.

And in the annealing process, after cooling to (Ms-150) degreeC or less, the time described in the "primary residence time" column of the hot rolling process of Table 2 at a temperature of (Ms-150) degrees C or less (8 to 37 s) holding (staying), and then performing a step of staying at 200°C or more and 440°C or less for 15 s or more.

In addition, in the case of manufacturing a cold-rolled steel sheet having no plating layer, in the annealing process, it is kept in a temperature range of 200°C or higher and 440°C or lower for 15 seconds or longer, and then quenched with water to room temperature (average cooling rate: about 50°C/ s) made it. In the case of manufacturing a cold-rolled steel sheet having a hot-dip galvanized layer on its surface, it was immersed in a plating bath and then rapidly cooled to room temperature using water (average cooling rate of about 50°C/s). In the case of manufacturing a cold-rolled steel sheet having an alloyed hot-dip galvanized layer on its surface, after performing an alloying treatment at an alloying temperature, it was rapidly cooled to room temperature with water (average cooling rate of about 50°C/s).

A test piece was sampled from the cold-rolled steel sheet, hot-dip-plated steel sheet, or alloyed hot-dip-plated steel sheet having no plating layer obtained as described above, and evaluated by the following method.

(i) tissue observation

The area ratio of each phase was evaluated by the following method. From the obtained steel sheet, cut out so that the cross section parallel to the rolling direction becomes the observation surface, the cross section is corroded with 1% nital, and enlarged by 2000 times with a scanning electron microscope, The inside of the area was photographed for 10 visual fields. t is the thickness (board thickness) of the steel plate. The ferrite phase is a structure in which corrosion marks or iron-based carbides are not observed in the mouth, and the tempered martensite phase is a structure in which a number of fine iron-based carbides and corrosion marks having an orientation in crystal grains are confirmed, and the martensite is not tempered. The trachea is a tissue in which no iron-based carbide is found in the mouth and is observed with white contrast under a scanning electron microscope. Since the grain boundary was also observed in white contrast, among the tissues observed in white contrast, linear tissues (an aspect ratio calculated as the length of the long axis/the length of the short axis is 10 or more) were excluded from the untempered martensite phase. Structures other than the ferrite phase, the tempered martensite phase, and the non-tempered martensite phase are shown in Table 3 as "other metal structures". In addition, in the column of "other metal structures" in Table 3, B is bainite, and residual γ is retained austenite.

Ferrite phase average particle diameter (in Table 3, described as ``ferrite average particle diameter''), ferrite phase area ratio (in Table 3, described as ``ferrite area ratio''), tempered martensite phase area ratio (in Table 3 (Referred to as ``tempered martensite area ratio''), and the interface length of the untempered martensite phase and the ferrite phase, and the interface length of the untempered martensite phase and the tempered martensite phase per unit area. The total value (described as "untempered martensite interface length" in Table 3) was obtained by image analysis of the observation result by the scanning electron microscope. Image analysis was performed using image analysis software (Image-Pro Plus ver. 7.0, manufactured by Nippon Ropa Co., Ltd.). The area ratio of the ferrite phase was obtained by extracting only the ferrite phase portion in each observation field, obtaining the area ratio occupied by the ferrite phase with respect to the observation field area, and averaging the value of the area ratio in 10 fields. Similarly, as for the area ratio of the tempered martensite phase, in each observation field, only the portion of the tempered martensite phase was extracted, and the area ratio occupied by the tempered martensite image was determined with respect to the observation field area, and the area ratio of 10 fields of view It was calculated|required by averaging the value of. In addition, the average particle diameter of the ferrite phase is obtained by obtaining the equivalent circle diameter corresponding to the area for each ferrite particle in each observation field, obtaining the average value of the circle equivalent diameter of the ferrite particles in the observation field, and this in each observation field. It was set as the circle-equivalent diameter of ferrite particles, and was calculated|required by averaging the value of the circle-equivalent diameter of the ferrite particles in 10 fields. ``Per unit area, the total length of the interface length of the untempered martensite phase and the ferrite phase, and the interface length of the non-tempered martensite phase and the tempered martensite phase" is tempered by image analysis in each observation field. The interface of the untempered martensite phase is determined, the total length of the interface of the non-tempered martensite phase and the ferrite phase present in the observation field, and the non-tempered martensite phase and the tempered martens present in the observation field The sum of the lengths of the interface of the zite phase is obtained, and the total value of both is divided by the observation field area, and the interfacial length of the untempered martensite and ferrite phases per unit area in each observation field, and the untempered marten It is the total value of the interface length of the zite phase and the tempered martensite phase, and the interface length of the untempered martensite phase and the ferrite phase per unit area of 10 fields of view, and the untempered martensite phase and the tempered martensite phase. It was calculated|required by averaging the total value of the interface length.

(ii) tensile test

A JIS No.5 tensile test piece was produced from the obtained steel sheet in a direction perpendicular to the rolling direction, and a tensile test in accordance with the provisions of JIS Z 2241 (2011) was performed 5 times, and the average yield strength (YS) and tensile strength (TS) , The total elongation (El) was determined. The crosshead speed in the tensile test was 10 mm/min. In Table 3, tensile strength: 980 MPa or more is a mechanical property required by the present invention.

(iii) bending evaluation

From the obtained steel sheet, a coil having a width of 100 mm with a direction parallel to the rolling direction as the bending test axis direction was roll-formed with R/t (R: bending radius, t: plate thickness) of 1.0 and 1.4. Thereafter, the test piece was visually observed to examine the presence or absence of cracks. In Table 3, the investigation result of the crack after performing the roll forming of R/t = 1.0 is "bending property at R/t = 1.0", and the investigation result of the crack after performing the roll forming of R/t = 1.4 is It is "bending property at R/t = 1.4". When a crack was confirmed, it was set as "○", and when a crack was not confirmed, it was set as "x", and if a crack did not occur at R/t of 1.4, it was set as pass (○).

(iv) Plating evaluation

Ten samples were taken from the obtained steel sheet for evaluation with a width of 800 mm and a length of 500 mm, and the presence or absence of subplating on the surface of the steel sheet was observed with the naked eye and a 10-fold loupe. In the case where no non-plating (region in which the plating was not formed) was observed, it was set to "○", and when the non-plating was confirmed, it was set to "x".

Table 3 shows the results obtained above.

In all of the examples of the present invention, it can be seen that the tensile strength TS: 980 MPa or more and good bendability were obtained. Moreover, in the example of this invention satisfying a specific condition, the plating property was also favorable. In addition, in these examples of the present invention, the hot-dip plated steel sheet having a hot-dip plated layer on the surface was shown, but the cold-rolled steel sheet of the present invention has a good bendability with tensile strength TS: 980 MPa or more, similar to the hot-dip plated steel sheet of the present invention It can be said that, and both of the cold-rolled steel sheet and hot-dip plated steel sheet of this invention can be used suitably as an automobile part, for example.

On the other hand, the comparative examples outside the scope of the present invention did not reach a tensile strength of 980 MPa, or a good one was not obtained in the bendability evaluation.

Claims (9)

By mass%,
C: 0.07% or more and 0.14% or less, Si: 0.65% or more and 1.65% or less,
Mn: 1.8% or more and 2.6% or less, P: 0.05% or less,
S: 0.005% or less, Al: 0.08% or less,
N: 0.0060% or less, Ti: 0.005% or more and 0.030% or less,
B: 0.0002% or more and 0.0030% or less,
Cr: 0.01% or more and 0.40% or less and Mo: 0.01% or more and 0.50% or less, containing one or two selected from, satisfying the following formula (1), and having a component composition consisting of Fe and unavoidable impurities. ,
Ferrite phase average particle diameter of 1.5 μm or less, ferrite phase area ratio of 2% or more and 15% or less, tempered martensite phase area ratio of 75% or more and 96% or less, per unit area, non-tempered martensite phase and ferrite A high-strength steel sheet in which the total value of the interface length of the phase and the interface length of the untempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/m 2 or more and 5.0 × 10 ¹¹ µm/m 2 or less.
[Equation 1]

Here, [%M] (M = Cr, Mo, C) represents the content of each element in mass%, respectively. The method of claim 1,
In addition to the above component composition, in addition to the mass%,
V: 0.001% or more and 0.3% or less,
Cu: 0.001% or more and 0.1% or less,
And Ni: a high-strength steel sheet containing one or two or more selected from 0.001% or more and 0.2% or less. The method according to claim 1 or 2,
High-strength steel plate with hot-dip plated layer on its surface The method of claim 3,
The hot-dip plated layer, by mass%, Fe: 5.0% or more and 20.0% or less, Al: 0.001% or more and 1.0% or less, and Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, A high-strength steel sheet containing one or two or more selected from Li, Ti, Be, Bi, and REM in a total of 0% or more and 3.5% or less, and the balance is composed of Zn and unavoidable impurities. The method of claim 3,
The high-strength steel sheet in which the hot-dip plated layer is an alloyed hot-dip plated layer. The steel material having the component composition according to claim 1 or 2 is heated at 1100° C. or higher and 1300° C. or lower, hot-rolled at a finish rolling of 820° C. or higher, and cooling is started within 3 seconds after the finish rolling to finish. After cooling at an average cooling rate of 30° C./s or more and less than 80° C./s from the rolling temperature to 700° C., cooling at an average cooling rate of 10° C./s or less to the coiling temperature, and winding at a coiling temperature of 580° C. or more and 680° C. or less. Hot rolling process,
After the hot rolling step, a cold rolling step of performing cold rolling, and
After the cold rolling, more than 500 ℃ (Ac ₃ - 120) with an average heating rate is more than 4.5 ℃ / s in the temperature range of less than ℃ and heated to a - - (10 Ac ₃₎ ℃ or less than (Ac ₃ 50) ℃ , After cooling to (Ms-150) ℃ or less, and having an annealing process of staying for 15 seconds or more in a temperature range of 200 ℃ to 440 ℃,
Ferrite phase average particle diameter of 1.5 μm or less, ferrite phase area ratio of 2% or more and 15% or less, tempered martensite phase area ratio of 75% or more and 96% or less, per unit area, non-tempered martensite phase and ferrite A method for producing a high-strength steel sheet in which the total value of the interface length of the phase and the interface length of the non-tempered martensite phase and the tempered martensite phase is 6.3 × 10 ⁸ µm/m 2 or more and 5.0 × 10 ¹¹ µm/m 2 or less. The method of claim 6,
A method for producing a high-strength steel sheet having a hot-dip plating process of hot-dip plating by heating again to 450°C or more and 600°C or less after the annealing process. The method of claim 7,
In the hot-dip plating process, a method for producing a high-strength steel sheet which is further subjected to alloying treatment after hot-dip plating treatment. The method of claim 4,
The high-strength steel sheet in which the hot-dip plated layer is an alloyed hot-dip plated layer.